Hot-rolled steel sheet having good cold workability and excellent hardness after working

ABSTRACT

This hot-rolled steel sheet has a thickness of 3-20 mm and contains specific amounts of C, Si, Mn, P, S, Al and N with the balance made up of iron and unavoidable impurities. The contents of solid-solved N, C and N are within specific ranges, and bainitic ferrite having a specific average crystal grain size and pearlite have specific area occupancies in the structure, with the balance occupied by polygonal ferrite. This hot-rolled steel sheet has a specific hardness distribution in the thickness direction.

TECHNICAL FIELD

The present invention relates to a hot-rolled steel sheet which shows satisfactory cold workability (heavy cold workability) when cold-worked so as to locally undergo an extremely large deformation strain and which shows given hardness after the working.

BACKGROUND ART

In recent years, from the standpoint of environmental protection, lighter weight, i.e., higher strength, of steel materials for use in various parts for automotive, for example, transmission parts such as gear, and casings, is increasingly required with the purpose of enhancing the fuel efficiency of automobiles. To meet this requirement for lighter weight and higher strength, a steel material prepared by hot-forging a steel bar (hot-forged material) has been used as a commonly-employed steel material. In addition, in order to reduce CO₂ emission in the process of producing parts, a requirement for cold forging of parts such as a gear, which had been heretofore worked by hot forging, is also more and more increasing.

Cold working (cold forging) is advantageous in that the productivity is high compared with hot working and warm working and moreover, both the dimensional accuracy and the steel material yield are good. The problem occurring in the case of producing parts by the cold working is that a steel material having high strength, i.e., high deformation resistance, must be necessarily used so as to ensure that the strength of cold-worked parts is equal to or more than a predetermined value expected. However, a higher deformation resistance of a steel material used leads to a shortening of the life of a metal mold for cold working.

In the field of transmission parts, investigations are being made on the production of parts from steel sheets, in place of production of forged products from steel bars (hot forging, cold forging, etc.), for the purpose of weight and cost reductions in parts. However, since transmission parts have complicated shapes, the parts produced form steel sheets by cold working (press forming, forging, etc.) have a drawback in that these parts locally have portions having an extremely large deformation strain (about 2 or larger in terms of true strain amount) and local cracking is apt to occur therein.

Therefore, conventionally, a method where a steel material is cold-forged into a predetermined shape and then subjected to a heat treatment such as quenching-tempering to produce a high-strength part assured of predetermined strength (hardness) is sometimes conducted. However, the heat treatment after cold forging inevitably causes a change in part dimension and therefore, it must be secondarily corrected by machining such as cutting. A possible solution to omit the step of heat treatment or subsequent working has been demanded.

In order to solve the problems above, for example, it is disclosed that when the progress of natural aging is restrained by using solute C in a low-carbon steel to ensure a predetermined amount of age hardening due to strain aging, a wire rod/steel bar for cold forging, excellent in the strain aging property, can be obtained (see, Patent Document 1).

However, in this technique, strain aging is controlled only by the solute C amount, and a steel material having sufficient cold workability as well as predetermined surface quality and hardness/strength after working can be hardly obtained.

Then, the present applicant had made various studies by focusing the attention on the difference of the effects of solute C and solute N contained in a steel material on the deformation resistance and static strain aging. As a result, it was found that when the amounts of these solute elements are appropriately controlled, a steel material for mechanical structure exerting good cold workability during working and exhibiting predetermined hardness (strength) after cold working (cold forging) can be obtained. The present applicant has already filed a patent application based on this finding (see, Patent Document 2).

This steel material realizes both cold workability and higher hardness (higher strength) after working but is a hot-forged material, similarly to the wire rod/steel bar described in Patent Document 1, and the production cost is disadvantageously high. In order to more reduce the production cost, studies are also being made to produce automobile parts by cold working by using a hot-rolled steel sheet in place of the conventional hot-forged material.

For example, a hot-rolled steel sheet for nitriding treatment, ensuring that high surface hardness and sufficient hardening depth are obtained after nitriding treatment, has been proposed (see, Patent Document 3).

However, this technique further requires a nitriding treatment after cold working and has a problem that a sufficient cost reduction cannot be realized.

In addition, a hot-rolled steel sheet having a composition containing C: 0.10% or less, Si: less than 0.01%, Mn: 1.5% or less, and Al: 0.20% or less, containing (Ti+Nb)/2: from 0.05 to 0.50%, and containing S: 0.005% or less, N: 0.005% or less, and O: 0.004% or less, with the total of S, N and O being 0.0100% or less, where the microstructure is a substantially ferrite single phase of 95% or more, has been proposed. This hot-rolled steel sheet is thought to be excellent in the dimensional accuracy of a finely blanked surface, ensure very high surface hardness of the blanked surface after working, and also be excellent in the resistance to red-scale defect (see, Patent Document 4).

However, this hot-rolled steel sheet where N is limited to a very low content as a harmful element utterly differs in the technical idea from the hot-rolled steel sheet according to the present invention where N is positively utilized.

RELATED ART Patent Document

Patent Document 1: JP-A-H10-306345

Patent Document 2: JP-A-2009-228125

Patent Document 3: JP-A-2007-162138

Patent Document 4: JP-A-2004-137607

SUMMARY OF THE INVENTION Problems that the Invention is to Solve

An object of the present invention, which has been achieved under these circumstances, is to provide a hot-rolled steel sheet which shows satisfactory cold workability (heavy cold workability) when cold-worked so as to undergo an exceedingly high degree of deformation strains and which shows given hardness after the working.

Means for Solving the Problems

In a hot-rolled steel sheet excellent in heavy cold workability and surface hardness after working in the first invention of the present invention,

a sheet thickness is from 3 to 20 mm;

a component composition comprises, in mass %,

C: more than 0% and 0.3% or less,

Si: more than 0% and 0.5% or less,

Mn: from 0.2 to 1%,

P: more than 0% and 0.05% or less,

S: more than 0% and 0.05% or less,

Al: from 0.01 to 0.1%, and

N: from 0.008 to 0.025%,

with the remainder being iron and unavoidable impurities, wherein

solute N: 0.007% or more and

the contents of C and N satisfy the relationship 10C+N≦3.0;

a microstructure comprises, in terms of area ratio relative to an entire microstructure, bainitic ferrite: 5% or more, pearlite: less than 20%, and remainder: polygonal ferrite;

an average grain size of the bainitic ferrite is in a range of from 3 to 50 μm; and

(Hv_(max)−Hv_(min))/Hv_(min)≦0.3, wherein, in a hardness distribution in a thickness direction, Hv_(max) and Hv_(min) are respectively the maximal value and the minimal value of Vickers hardness values of three portions that are a surface portion, a t/4 portion where t is the sheet thickness, and a central portion.

In the hot-rolled steel sheet excellent in heavy cold workability and surface hardness after working in the second invention of the present invention according to the first invention, the component composition further comprises, in mass %:

at least one member selected from the group consisting of Cr: more than 0% and 2% or less and Mo: more than 0% and 2% or less.

In the hot-rolled steel sheet excellent in heavy cold workability and surface hardness after working in the third invention of the present invention according to the first or second invention, the component composition further comprises, in mass %:

at least one member selected from the group consisting of Ti: more than 0% and 0.2% or less, Nb: more than 0% and 0.2% or less, and V: more than 0% and 0.2% or less.

In the hot-rolled steel sheet excellent in heavy cold workability and surface hardness after working in the fourth invention of the present invention according to any one of the first to third inventions, the component composition further comprises, in mass %.

B: more than 0% and 0.005% or less.

In the hot-rolled steel sheet excellent in heavy cold workability and surface hardness after working in the fifth invention of the present invention according to any one of the first to fourth inventions, the component composition further comprises, in mass %:

at least one member selected from the group consisting of Cu: more than 0% and 5% or less, Ni: more than 0% and 5% or less, and Co: more than 0% and 5% or less.

In the hot-rolled steel sheet excellent in heavy cold workability and surface hardness after working in the sixth invention of the present invention according to any one of the first to fifth inventions, the component composition further comprises, in mass %:

at least one member selected from the group consisting of Ca: more than 0% and 0.05% or less, REM: more than 0% and 0.05% or less, Mg: more than 0% and 0.02% or less, Li: more than 0% and 0.02% or less, Pb: more than 0% and 0.5% or less, and Bi: more than 0% and 0.5% or less.

Effects of the Invention

According to the present invention, the hot-rolled steel sheet has a microstructure which mainly contains bainitic ferrite having a given average grain size and polygonal ferrite and in which a solute N amount has been ensured and the content of C and the content of N satisfy a given relationship. Because of this, the hot-rolled steel sheet shows reduced deformation resistance during cold working, thereby prolonging the life of the die. Furthermore, since this hot-rolled steel sheet has a hardness distribution in a thickness direction regulated so as to be in a given range, this hot-rolled steel sheet is less apt to suffer local cracking even in cold working which may locally cause an extremely large deformation strain, and the parts obtained therefrom by working can have given hardness.

BRIEF DESCRIPTION OF THE DRAWING

FIG. 1 is a view which diagrammatically shows the configuration of the wedge type compression tester used in the Examples for evaluating heavy cold workability.

MODES FOR CARRYING OUT THE INVENTION

The hot-rolled steel sheet in the present invention (hereinafter referred to also as “steel sheet of the present invention” or merely as “steel sheet”) is explained below in more detail. The steel sheet of the present invention has the features of ensuring a solute N amount and having C and N contents which satisfy a given relationship, in common with the hot-forged material described in Patent Document 2. However, the steel sheet of the present invention differs from the hot-forged material described in Patent Document 2 in that the steel sheet of the present invention has a C content permissible to a relatively high value and has a bainitic ferrite/polygonal ferrite/pearlite multiphase microstructure, that the bainitic ferrite grains are refined, and that the steel sheet has a hardness distribution in a thickness direction regulated so as to be in a given range.

[Thickness of Steel Sheet of the Present Invention: From 3 to 20 mm]

First, a steel sheet having a thickness of 3 to 20 mm is targeted by the steel sheet of the present invention. If the sheet thickness is less than 3 mm, the rigidity as a structure cannot be ensured. On the other hand, if the sheet thickness exceeds 20 mm, the microstructure specified in the present invention can be hardly achieved, and the desired effects cannot be obtained. The sheet thickness is preferably from 4 to 19 mm.

Next, the component composition constituting the steel sheet of the present invention is described. In the following, the units of chemical components are all mass %.

[Component Composition of Steel Sheet of the Present Invention]

<C: More than 0% and 0.3% or Less>

C is an element greatly affecting the formation of the microstructure of the steel sheet and although the microstructure is a bainitic ferrite-polygonal ferrite-pearlite multi-phase microstructure, in order to form a bainitic ferrite-polygonal ferrite-based microstructure with as little pearlite as possible, the content of this element needs to be limited. If C is contained too much, the pearlite fraction in the steel sheet microstructure is increased, leaving a fear that the deformation resistance becomes excessive due to work hardening of pearlite. Therefore, the C content in the steel sheet is limited to 0.3% or less, preferably 0.25% or less, more preferably 0.2% or less, and still more preferably 0.15% or less. However, if the C content is too small, deoxidation during melting of steel is difficult to be achieved and at the same time, the strength and hardness after cold working can be hardly satisfied. Therefore, it is preferably 0.0005% or more, more preferably 0.0008% or more and still more preferably 0.001% or more.

<Si: More than 0% and 0.5% or Less>

Si forms a solid solution in steel to thereby increase the deformation resistance of the steel sheet and thus, is an element needed to be reduced as much as possible. Therefore, in order to suppress the increase in deformation resistance, the Si content in the steel sheet is limited to 0.5% or less, preferably 0.45% or less, more preferably 0.4% or less, and still more preferably 0.3% or less. However, if the Si content is extremely small, deoxidation during melting is difficult to be achieved and at the same time, the strength and hardness after cold working can be hardly satisfied. Therefore, it is preferably 0.005% or more, more preferably 0.008% or more, and still more preferably 0.01% or more.

<Mn: From 0.2 to 1%>

Mn is an element exerting deoxidation and desulfurization actions in the process of steel making. Furthermore, when the N content in the steel material is increased, cracking is readily generated due to dynamic strain aging by heat generation during working, but, on the other hand, Mn has an effect of enhancing the workability on this occasion and inhibiting cracking. In order to effectively bring out these actions, the Mn content in the steel sheet is 0.2% or more, preferably 0.22% or more and more preferably 0.25% or more. However, if the Mn content is too large, the deformation resistance becomes excessive, and segregation occurs to produce a heterogeneity in the microstructure. Therefore, it is 1% or less, preferably 0.98% or less and more preferably 0.95% or less.

<P: More than 0% and 0.05% or Less>

P is an impurity element unavoidably contained in the steel. It is an element that, if contained in ferrite, segregates at a ferrite grain boundary to deteriorate cold workability and contributes to solid-solution hardening of ferrite and thereby gives rise to an increase in the deformation resistance. Therefore, the P content is preferably reduced as much as possible in view of cold workability, but if excessively reduced, the steel making cost increases. Therefore, in consideration of process capability, the content is 0.05% or less and preferably 0.03% or less.

<S: More than 0% and 0.05% or Less>

S is also an unavoidable impurity, similarly to P, and is an element precipitating as FeS at a grain boundary in a film form to deteriorate the workability. In addition, this also has an action of causing hot shortness. In the present invention, from the standpoint of enhancing the deformation performance, the S content is 0.05% or less and preferably 0.03% or less. However, reduction of the S content to 0 is difficult in industry. Since S has an effect of enhancing the machinability, in view of machinability enhancement, it is recommended to be contained in an amount of preferably 0.002% or more and more preferably 0.006% or more.

<Al: From 0.01 to 0.1%>

Al is an element effective for deoxidation in the process of steel making. In order to obtain this deoxidation effect, the Al content in the steel sheet is 0.01% or more, preferably 0.015% or more and more preferably 0.02% or more. However, if the Al content is excessively large, toughness is reduced and cracking readily occurs. Therefore, the content is 0.1% or less, preferably 0.09% or less and more preferably 0.08 mass % or less.

<N: From 0.008 to 0.025%>

N is an important element for obtaining predetermined strength by static strain aging after working. Therefore, the N content in the steel sheet is 0.008% or more, preferably 0.0085% or more and more preferably 0.009% or more. However, if the N content is excessively large, the effect of dynamic strain aging during working, in addition to static strain aging, becomes significant, and thus the deformation resistance is increased, which is unsuitable. Therefore, the content is 0.025% or less, preferably 0.023% or less and more preferably 0.02% or less.

<Solute N: 0.007% or More>

Furthermore, by ensuring a content of solute N (hereinafter referred to as “solute N amount”) in the steel sheet, the static strain aging can be promoted without increasing the deformation resistance very much. For ensuring a required strength after cold working, a solute N amount of 0.007% or more is necessary. However, in a case where the solute N amount is too high, not only the cold workability is deteriorated, but also an amount of the solute N fixed to working strains is increased and this is prone to result in a hardness distribution along the thickness direction of the hot-rolled sheet. This hardness distribution in a thickness direction cannot be eliminated even when the annealing conditions which will be described later are applied, and this hot-rolled steel sheet is prone to crack when subjected to working which may cause an extremely large local deformation strain. Consequently, the solute N amount is preferably 0.03% or less. Since the content of N in the steel material is 0.025% or less, the case where the solute N amount is 0.025% or more is not substantially occurred.

Here, the solute N amount in the present invention is an amount determined by subtracting the amount of total N compounds from the total N amount in the steel sheet in conformity with JIS G 1228. An example of the practical method for measuring the solute N amount is described below.

(a) Inert Gas Fusion Method-Thermal Conductivity Method (Measurement of Total N Amount)

A sample cut out from a test material is placed in a crucible and fused in an inert gas stream to extract N, and the extract is transferred to a thermal conductivity cell and measured for the change in thermal conductivity to determine the total N amount.

(b) Ammonia Distillative Separation and Indophenol Blue Absorptiometry (Measurement of Amount of Total N Compounds)

A sample cut out from a test material is dissolved in a 10% AA-type electrolytic solution and a constant current electrolysis is performed to measure the amount of total N compounds in the steel. The 10% AA-type electrolytic solution used is a nonaqueous solvent-type electrolytic solution composed of 10% acetone and 10% tetramethylammonium chloride, with the remainder being methanol, and is a solution not forming a passive film on the steel surface.

About 0.5 g of the sample of the test material is dissolved in the 10% AA-type electrolytic solution, and the insoluble residue (N compounds) produced is filtered through a polycarbonate-made filter having a pore size of 0.1 μm. The obtained insoluble residue is decomposed by heating in sulfuric acid, potassium sulfate and pure copper-made chips, and the decomposition product is combined with the filtrate. The resulting solution is made alkaline with sodium hydroxide and then subjected to steam distillation, and the distilled ammonia is absorbed by diluted sulfuric acid. Furthermore, phenol, sodium hypochlorite and sodium pentacyanonitrosylferrate(III) are added to produce a blue complex, and the absorbance thereof is measured by using an absorptiometer to determine the amount of total N compounds.

The solute N amount can be determined by subtracting the amount of total N compounds determined by the method (b) from the total N amount determined by the method (a).

<Contents of C and N Satisfy the Relationship of 10C+N≦3.0>

In the steel material of the present invention, solute C greatly increases the deformation resistance and does not so much contribute to static strain aging and, on the other hand, the solute N can promote the static strain aging without raising the deformation resistance very much and therefore, has an action of allowing for an increase in the hardness after working. Therefore, in the steel material of the present invention, in order to increase the hardness after working without raising the deformation resistance during working very much, the C content and the N content must satisfy the relationship of 10C+N≦3.0. It is preferably 0.009≦10C+N≦2.8, more preferably 0.01≦10C+N≦2.5 and still more preferably 0.01≦10C+N≦2.0. From the standpoint of refining the grain in the hot-rolled steel sheet and ensuring formability of the steel sheet, the C content and the solute C amount are needed to some extent, but if 10C+N>3.0, the amounts of C and/or N are too large, and the deformation resistance becomes excessive. In the inequality above, the coefficient of the C content is set to be 10 times the coefficient of the N content by taking into account the fact that even when the contents are the same, the degree of increase in the strength and deformation resistance in the hot-rolled steel sheet of the present invention, which is attributable to the solute C, is about one digit (10 times) larger than that attributable to the solute N.

The steel of the present invention fundamentally contains the above-described components, with the remainder being iron and unavoidable impurities, but in addition, the following allowable components may be added, as long as the action of the present invention is not impaired.

<At Least One Member Selected from the Group Consisting of Cr: More than 0% and 2% or Less and Mo: More than 0% and 2% or Less>

Cr is an element having an action of increasing the grain boundary strength and thereby enhancing the deformation performance of the steel. In order to effectively bring out such an action, Cr is preferably contained in an amount of 0.2% or more, but if Cr is contained too much, the deformation resistance may be increased to reduce the cold workability. Therefore, it is recommended that the content thereof is 2% or less, furthermore 1.5% or less, and in particular 1% or less.

Mo is an element having an action of increasing the hardness of the steel material after working and the deformation performance. In order to effectively bring out such an action, Mo is preferably contained in an amount of 0.04% or more, more preferably 0.08% or more. However, if Mo is contained too much, the cold workability may be deteriorated. Therefore, it is recommended that the content thereof is 2% or less, furthermore 1.5% or less, and in particular 1% or less.

<At Least One Member Selected from the Group Consisting of Ti: More than 0% and 0.2% or Less, Nb: More than 0% and 0.2% or Less and V: More than 0% and 0.2% or Less>

These elements have a high affinity for N and are elements fulfilling the role of forming N compounds by coexisting with N, refining the grain of steel, enhancing the toughness of a processed product obtained after cold working, and also enhancing the cracking resistance. Even if each element is contained in an amount over the upper limit value, an effect of improving the property is not obtained. Therefore, it is recommended that the content of each element is 0.2% or less, furthermore from 0.001 to 0.15% and in particular from 0.002 to 0.1%.

<B: More than 0% and 0.005% or Less>

Similarly to Ti, Nb and V above, B has a high affinity for N and is an element fulfilling the role of forming a N compound by coexisting with N, refining the grain of steel, enhancing the toughness of a processed product obtained after cold working, and also enhancing the cracking resistance. Therefore, in the case where the steel sheet of the present invention contains B, a predetermined solute N amount can be ensured to enhance the strength after cold working. For this reason, it is recommended that the content thereof is 0.005% or less, furthermore from 0.0001 to 0.0035% and in particular from 0.0002 to 0.002%.

<At Least One Member Selected from the Group Consisting of Cu: More than 0% and 5% or Less, Ni: More than 0% and 5% or Less and Co: More than 0% and 5% or Less>

All of these elements have an action of hardening the steel material by stain aging and are elements effective for enhancing the post-working strength. In order to effectively bring out such an action, each of these elements is preferably contained in an amount of 0.1% or more and furthermore 0.3% or more. However, if the content of each of these elements is too much large, the effect of hardening the steel material by stain aging and furthermore the effect of enhancing the post-working strength may be saturated, or the cracking may be promoted. Therefore, it is recommended that each of them is 5% or less, furthermore 4% or less and in particular 3% or less.

<At Least One Member Selected from the Group Consisting of Ca: 0.05% or Less (Exclusive of 0%), REM: 0.05% or Less (Exclusive of 0%), Mg: 0.02% or Less (Exclusive of 0%), Li: 0.02% or Less (Exclusive of 0%), Pb: 0.5% or Less (Exclusive of 0%), and Bi: 0.5% or Less (Exclusive of 0%)>

Ca is an element spheroidizing a sulfide compound-based inclusion such as MnS to thereby enhance the deformation performance of steel and at the same time, contributing to improvement of the machinability. In order to effectively bring out such an action, Ca is preferably contained in an amount of 0.0005% or more and furthermore 0.001% or more. Even if contained too much, the effect thereof is saturated and an effect consistent with the content cannot be expected. Therefore, 0.05% or less, furthermore 0.03% or less and in particular 0.01% or less are recommended.

REM is, similarly to Ca, an element spheroidizing a sulfide compound-based inclusion such as MnS to thereby enhance the deformation performance of steel and at the same time, contributing to improvement of the machinability. In order to effectively bring out such an action, REM is preferably contained in an amount of 0.0005% or more and furthermore 0.001% or more. Even if contained too much, the effect thereof is saturated and an effect consistent with the content cannot be expected. Therefore, 0.05% or less, furthermore 0.03% or less and in particular 0.01% or less are recommended.

The “REM” as used in the present invention means to include lanthanoid elements (15 elements from La to Lu) as well as Sc (scandium) and Y (yttrium). Among these elements, it is preferable to contain at least one element selected from the group consisting of La, Ce and Y, and it is more preferable to contain La and/or Ce.

Mg is, similarly to Ca, an element spheroidizing a sulfide compound-based inclusion such as MnS to thereby enhance the deformation performance of steel and at the same time, contributing to improvement of the machinability. In order to effectively bring out such an action, Mg is preferably contained in an amount of 0.0002% or more and furthermore 0.0005% or more. Even if contained too much, the effect thereof is saturated and an effect consistent with the content cannot be expected. Therefore, 0.02% or less, furthermore 0.015% or less and in particular 0.01% or less are recommended.

Li is, similarly to Ca, an element spheroidizing a sulfide compound-based inclusion such as MnS to allow for enhancement of the deformation performance of steel and in addition, contributing to improvement of the machinability by lowering the melting point of an Al-based oxide and thereby making it harmless. In order to effectively bring out such an action, Li is preferably contained in an amount of 0.0002% or more and furthermore 0.0005% or more. Even if contained too much, the effect thereof is saturated and an effect consistent with the content cannot be expected. Therefore, 0.02% or less, furthermore 0.015% or less and in particular 0.01% or less are recommended.

Pb is an element effective for enhancing the machinability. In order to effectively bring out such an action, Pb is preferably contained in an amount of 0.005% or more and furthermore 0.01% or more. However, if contained too much, there arises a problem with production such as formation of a roll mark. Therefore, 0.5% or less, furthermore 0.4% or less and in particular 0.3% or less are recommended.

Bi is, similarly to Pb, an element effective for enhancing the machinability. In order to effectively bring out such an action, Bi is preferably contained in an amount of 0.005% or more and furthermore 0.01% or more. Even if contained too much, the effect of enhancing the machinability is saturated. Therefore, 0.5 mass % or less, furthermore 0.4% or less and in particular 0.3% or less are recommended.

The microstructure characterizing the steel sheet of the present invention is described below.

[Microstructure of Steel Sheet of the Present Invention]

Although the steel sheet of the present invention is based on a steel having a bainitic ferrite/polygonal ferrite/pearlite multiphase microstructure as described above, this steel sheet is especially characterized in that the size of the bainitic ferrite grains has been regulated so as to be in a specific range and that the hardness distribution in a thickness direction has been regulated.

<Bainitic Ferrite: 5% or More, Pearlite: Less than 20%, and Remainder: Polygonal Ferrite>

The microstructure of the steel sheet of the present invention is constituted of a multiphase microstructure composed of bainitic ferrite, polygonal ferrite, and pearlite. Bainitic ferrite not only has, during cold working, the function of enhancing the workability but also has, after the wording, the function of increasing the hardness, and suppressing the generation of stretcher strain marks. From the standpoint of making the bainitic ferrite effectively perform these functions, the area ratio thereof is 5% or more, preferably 10% or more, more preferably 15% or more. The upper limit of the area ratio of bainitic ferrite in the steel sheet of the present invention is substantially about 90%, preferably 85%, more preferably 80%. Meanwhile, in a case where pearlite is present in too large amount, the formability of the steel sheet is deteriorated. Consequently, the area ratio of pearlite is 20% or less, preferably 19% or less, more preferably 18% or less, especially preferably 15% or less. The lower limit of the area ratio of pearlite in the steel sheet of the present invention is substantially about 0.5%, preferably 1%. Although the remainder is polygonal ferrite, the area ratio of polygonal ferrite is preferably 5% or more.

In the microstructure of the steel sheet of the present invention, a cementite phase is present besides the microstructures described above. However, the area ratio thereof is as slight as about 1% at the most. Consequently, in this description, the area ratios of bainitic ferrite, polygonal ferrite, and pearlite were normalized and defined so that the total area ratio of these three phases is 100%.

<Average Grain Size of Bainitic Ferrite: In a Range of from 3 to 50 μm>

The average grain size of bainitic ferrite constituting the bainitic ferrite microstructure must be in a range of from 3 to 50 μm so as to enhance the workability of the steel sheet and satisfy the surface property after working. If the bainitic ferrite particles are excessively fine, the deformation resistance becomes too high. Therefore, the average grain size thereof is 3 μm or more, preferably 4 μm or more and more preferably 5 μm or more. On the other hand, if the bainitic ferrite is excessively coarsened, the surface property after working is deteriorated and in addition, toughness, fatigue property, etc. are reduced. Therefore, the average grain size thereof is 50 μm or less, preferably 45 μm or less and more preferably 40 μm or less.

<Hardness Distribution in Thickness Direction: (Hv_(max)−Hv_(min))/Hv_(min) Regulated to 0.3 or Less, Wherein Hv_(max) and Hv_(min) are Respectively the Maximal Value and the Minimal Value of Vickers Hardness Values of Three Portions that are a Surface Portion, a t/4 Portion where t is the Sheet Thickness, and a Central Portion>

Since transmission parts have complicated shapes, press forming or forging results in regions which locally have an extremely large deformation strain (corresponding to a true strain ε of about 2 or larger). In the case of steel sheets having a large hardness distribution in a thickness direction (strength distribution, stress distribution), such large local deformation strains undesirably result in uneven plastic deformations. In lowly worked regions, i.e., regions having a small deformation strain (less than about 2 in terms of ε), the hardness distribution in a thickness direction exerts little influence and arouses no problem. However, in regions having a large strain amount (about 2 or larger in terms of ε), the hardness distribution in a thickness direction undesirably results in local cracking. In order to prevent local cracking from occurring even in such regions having a strain amount as extremely large as about 2 in terms of ε, the value of (Hv_(max)−Hv_(min))/Hv_(min), wherein Hv_(max) and Hv_(min) are respectively the maximal value and the minimal value of Vickers hardness values of three portions that are a surface portion, a t/4 portion where t is the sheet thickness, and a central portion, is regulated, as a hardness distribution in a thickness direction, to 0.3 or less, preferably 0.2 or less, more preferably 0.15 or less.

The mechanism by which a hardness distribution in a thickness direction arises in a conventional hot-rolled steel sheet is presumed to be as follows. In a hot-rolled steel sheet having a large sheet thickness, examples of the causes of the occurrence of a hardness distribution in a thickness direction include a difference in the degree of working between each surface portion and the central portion and a difference in working temperature (including the heat generated by the working) between each surface portion and the central portion, these differences unavoidably occurring during the hot working. Furthermore, phase transformations, generation of residual stress, and the like which occur during coil cooling also exert influences. In the present invention, since the alloying components thereof include solute N in a large amount, the fixing of N to regions having a large working strain occurs to undesirably increase the hardness of such regions having a large working strain, and this increase in hardness also exerts an influence. A plurality of such complicated factors cause a hardness distribution in a thickness direction, which is prone to result in thickness-direction unevenness in strength.

The steel sheet of the present invention can hence be obtained by subjecting a sheet which has just been hot-rolled to batch annealing under the recommended conditions which will be described later to thereby reduce the hardness distribution in a thickness direction.

[Method for Measuring Area Ratio of Each Phase]

As for the area ratio of the each phase above, each test steel sheet is subjected to Nital etching, and five visual fields are photographed by a scanning electron microscope (SEM, magnification: 1,000 times), and as a result, respective percentages of bainitic ferrite, polygonal ferrite and pearlite can be determined by a point counting method.

Here, the bainitic ferrite is defined as a ferrite particle existing in the bainite (collectively referring to upper bainite and lower bainite) microstructure in which the grain is in an axially extended shape (see, Tadashi Furuhara, “Current Opinion on Definition of Bainite Structure in Steels”, Netsu Shori, Vol. 50, No. 1, February 2010, pp. 22-27) and an aspect ratio (major axis/minor axis ratio) is 2 or more. In addition, the polygonal ferrite is defined as a ferrite particle in which the grain is in an equiaxed shape and an aspect ratio (major axis/minor axis ratio) is less than 2.

[Method for Measuring Average Grain Size]

The average grain size of the bainitic ferrite above can be measured as follows. That is, the grain sizes of bainitic ferrite present at three portions, i.e., an outermost layer portion, a portion at ¼ of the sheet thickness, and a central portion in the sheet thickness direction, are measured. As to the grain size of one bainitic ferrite particle, the side surface part in the rolling direction at each measurement portion is subjected to Nital etching, five visual fields of the corresponding region are photographed by a scanning electron microscope (SEM; magnification: 1,000 times), and the diameter including the center of gravity of the bainitic ferrite grain is determined by image analysis and defined as the average grain size.

[Method for Determining Hardness Distribution in Thickness Direction]

A thickness-direction cross-section parallel with the rolling direction of the hot-rolled sheet was examined for Vickers hardness (Hv) with respect to each of a surface portion (located at a depth of 400 μm form a sheet surface), a portion at ¼ of the sheet thickness, and a central portion in the sheet shickness direction, using a micro-Vickers hardness tester under the conditions of a load of 50 g and the number of measurements of 5 times. An average for the five measurements was taken as the Vickers hardness of each portion.

Of these Vickers hardness values for the three portions, the maximal value Hv_(max) and the minimal value Hv_(min) were determined to calculate (Hv_(max)−Hv_(min))/Hv_(min).

A preferable production method for obtaining the above-described steel sheet of the present invention is described below.

[Preferable Method for Producing Steel Sheet of the Present Invention]

The production of the steel sheet of the present invention may be conducted according to any method as long as it is a method capable of forming a raw material steel having the above-described chemical composition into a desired thickness. For example, it can be conducted by a method in which, under following conditions, a molten steel having the above-described component composition is prepared in a converter, subjected to ingot making or continuous casting to form a slab, and then rolled into a hot-rolled steel sheet having a desired thickness.

[Preparation of Molten Steel]

The N content in the molten steel can be adjusted by adding a N compound-containing raw material to the molten steel and/or controlling the atmosphere of the converter to a N₂ atmosphere, during melting in the converter.

[Heating]

Heating before hot rolling is performed at 1,100 to 1,300° C. In this heating, a high-temperature heating condition is necessary so as to produce no N compound and dissolve as much N as possible in solid. As for the heating temperature, preferable lower limit is 1,100° C. and more preferable lower limit is 1,150° C. On the other hand, a temperature more than 1,300° C. is operationally difficult.

[Hot Rolling]

Hot rolling is performed such that the finish rolling temperature is 880° C. or higher. If the finish rolling temperature is too low, ferrite transformation takes place at a high temperature, leading to coarsening of the precipitated carbide in ferrite (collectively referring to bainitic ferrite and polygonal ferrite), and the fatigue strength is deteriorated. Therefore, a finish rolling temperature not less than a certain level is necessary. The finish rolling temperature is more preferably 900° C. or more so as to coarsen the austenite particle and thereby increase the grain size of bainitic ferrite to a certain extent. The upper limit of the finish rolling temperature is 1,000° C. because temperature ensuring is difficult.

The thickness of the hot-rolled steel sheet of the present invention is from 3 to 20 mm. In order to refine the bainitic ferrite grain and thereby control the average grain size thereof to fall in a predetermined grain size range, not only the rolling temperature must be controlled as above but also the final rolling reduction by tandem rolling in the finish rolling must be controlled to be 15% or more. Usually, in the finish rolling, tandem rollings of from 5 to 7 passes are conducted, where the pass schedule is set from the standpoint of controlling jamming of the sheet and the final rolling reduction is up to approximately from 12 to 13%. The final rolling reduction is preferably 16% or more and more preferably 17% or more. As the final rolling reduction is higher, e.g., 20% or 30%, the effect of more refining the grain is obtained, but in view of rolling control, the upper limit is specified to be about 30%.

[Rapid Cooling after Hot Rolling]

After the completion of the finish rolling, the sheet is rapidly cooled at a cooling rate (first cooling rate) of 20° C./s or more within 5 seconds and the rapid cooling is stopped at a temperature (rapid cooling stop temperature) of 550° C. or more and less than 650° C. This is performed so as to obtain a bainitic ferrite-polygonal ferrite-pearlite multi-phase microstructure having predetermined phase fractions. If the cooling rate (rapid cooling rate) is less than 20° C./s, pearlite transformation is promoted, and if the rapid cooling stop temperature is less than 550° C., bainite transformation is suppressed. In both cases, a bainitic ferrite-polygonal ferrite-pearlite steel having predetermined phase fractions can be hardly obtained, and the cold workability or surface quality after working is deteriorated. On the other hand, if the rapid cooling stop temperature is 650° C. or more, the precipitated carbide in ferrite is coarsened, and the fatigue strength is reduced. The rapid cooling stop temperature is preferably from 560 to 640° C. and more preferably from 580 to 620° C.

[Slow Cooling after Stopping of Rapid Cooling]

After stopping the rapid cooling, the sheet is slowly cooled by standing to cool or air cooling at a cooling rate (slow cooling rate) of 10° C./s or less for 5 to 20 seconds. Accordingly, the precipitated carbide in ferrite is appropriately refined while allowing polygonal ferrite formation to proceed sufficiently. If the cooling rate exceeds 10° C./s or the slow cooling time is less than 5 seconds, the amount of polygonal ferrite formed is insufficient, whereas if the slow cooling time exceeds 20 seconds, the precipitated carbide is not coarsened and the fatigue strength is deteriorated.

[Rapid Cooling and Coiling after Slow Cooling]

After the slow cooling, the sheet is again rapidly cooled at a cooling rate (second rapid cooling rate) of 20° C./s or more and coiled at 500 to 600° C. This is performed so as to form a bainitic ferrite+polygonal ferrite-based microstructure and thus ensure cold workability. If the cooling rate (second rapid cooling rate) is less than 20° C./s or the coiling temperature exceeds 600° C., the cold workability is deteriorated due to formation of a large amount of pearlite, whereas if it is less than 500° C., the amount of bainitic ferrite formed is insufficient and the surface quality after working is deteriorated.

[Batch Annealing after Hot Rolling]

After the hot rolling, the sheet which has just been hot-rolled (hot-rolled coil) is subjected to batch annealing under the following conditions in order to regulate the hardness distribution in a thickness direction so as to be in the given range.

Specifically, this batch annealing is conducted in an atmosphere having an H₂ concentration of from 15 to 20 vol % by heating the steel sheet from room temperature to a temperature which is 400° C. or higher but not higher than Ac1 and then holding the steel sheet for 1 hour or more and 15 hours or less, in order to suppress surface scale formation and decarbonization.

The holding temperature and the holding period vary depending on the thickness of the sheet which has just been hot-rolled and on the size of the coil, and are suitably selected in accordance with the required narrowness of the hardness distribution in a thickness direction, which corresponds to the degree of cold working to be required, and with the evenness of the internal temperature of the coil

This heat treatment serves not only to remove the residual stress generated during the hot-rolling, thereby softening the steel sheet and diminishing strains, but also to release the fixed N element to accelerate spheroidization of carbides. In addition, the heat treatment serves to dissolve fine lamellae in the austenite. The hardness distribution in a thickness direction is reduced thereby. After the batch annealing, the steel sheet is cooled to 600° C. at a rate of 10° C./h or less to thereby accelerate the spheroidization of carbides. Subsequently, the steel sheet is cooled from 600° C. to 400° C. at a rate of 15° C./h or less, for the purpose of evenly cooling the whole coil and thereby preventing coil collapse or the like to stabilize the shape. Thereafter, cooling from 400° C. may be performed at a higher cooling rate (e.g., about 50 to 100° C./h or higher) by water cooling, etc. so long as the coil can be cooled while maintaining an even temperature distribution within the coil.

In a case where the holding temperature in the batch annealing is lower than 400° C., those effects are insufficient. Meanwhile, in a case where the holding temperature exceeds the Ac1 point, the microstructure changes undesirably. The holding temperature is more preferably 450 to 650° C., especially preferably 500 to 600° C.

In a case where the holding period is less than 1 hour, those effects are insufficient. Meanwhile, holding periods exceeding 15 hours are undesirable because the effects cannot be enhanced any more, the production efficiency is impaired, and a surface scale is prone to generate. The holding period is more preferably 2 to 14 hours, especially preferably 3 to 12 hours.

The present invention is described in greater detail below by referring to Examples, but the present invention is by no means limited to the following Examples and may be carried out by appropriately making changes as long as they are in conformity to the gist described hereinabove and hereinafter, all of which are included in the technical scope of the present invention.

EXAMPLES

Steels having the component compositions shown in Table 1 below were produced by a vacuum melting method and cast into ingots having a thickness of 120 mm. These ingots were subjected to hot rolling and then batch annealing under the conditions shown in Table 2 and Table 3 below to produce hot-rolled steel sheets. In each test, the following conditions were used: the rate of cooling after completion of the finish rolling to a stop of rapid cooling was 20° C./s or higher, and the cooling after the stop of rapid cooling was slow cooling conducted for 5 to 20 seconds at a cooling rate of 10° C./s or less; and after the batch annealing, the steel sheet was cooled to 600° C. at a cooling rate of 10° C./h or less, subsequently cooled form 600° C. to 400° C. at a cooling rate of 15° C./h or less, and further cooled from 400° C. by water cooling.

The hot-rolled steel sheets thus obtained were each examined for solute N amount, area ratio of each phase in the microstructure of the steel sheet, average grain size of bainitic ferrite, and hardness distribution in a thickness direction by the measurement methods explained above in the section “MODES FOR CARRYING OUT THE INVENTION”.

Furthermore, the hot-rolled steel sheets were evaluated for heavy cold workability and hardness after working in the following manners.

(Evaluation of Heavy Cold Workability)

In order to evaluate cold workability which locally causes an extremely large deformation strain (heavy cold workability), the following test was conducted as a test in which a working strain of 4 or larger in terms of true strain was introduced into a surface portion of a test specimen. An 80-ton pressing test machine was used to perform a wedge type compression test in which a cylindrical test specimen and wedge type jigs were used, the configuration of the test being diagrammatically shown in FIG. 1 (the test specimen was compressed at a compression rate of 1 mm/sec to a reduction of 80% relative to the diameter thereof). Test specimens used were as follows. In the case of hot-rolled steel sheets having a thickness of 10 mm or larger, cylindrical test specimens having a diameter of 10 mm were cut out therefrom. In the case of hot-rolled steel sheets having a thickness less than 10 mm, cylindrical test specimens having a diameter equal to the sheet thickness were cut out therefrom.

Prior to this compression test, forging analysis software FORGE (manufactured by TRANSVALOR S.A.) was used to calculate a distribution of true strains within a test specimen at the time when the reduction in the compression test was 80%. It was thus ascertained that the true strain ε was 4 or larger in the position located at a depth of 100 μm from the surface portion compressed by the R part of the compression jig, among the surface portions of the test specimen.

The test specimen which had undergone the wedge type compression test was visually examined and evaluated for heavy cold workability in accordance with the following criteria. The case of “o” was regarded as acceptable.

-   -   ∘: no cracks occurred in the test specimen     -   Δ: minute cracks occurred in the surface of the test specimen     -   x: cracks occurred in the test specimen         (Evaluation of Hardness after Working)

Hardness after working was evaluated by measuring the Vickers hardness (Hv) of the center of the surface of that portion of the test specimen which had been compressed by the compression jig in the wedge type compression test, using a Vickers hardness tester under the conditions of a load of 500 g and the number of measurements of 5 times. An average thereof was taken as hardness after working. The steel sheets having a hardness after working of 250 Hv or more were regarded as acceptable.

The results of those measurements are shown in Tables 4 to 6 below.

TABLE 1 Kind of Components (mass %) [remainder: Fe and unavoidabe impurities] steel C Si Mn P S Al N 10 C + N Others a 0.05 0.02 0.40 0.007 0.001 0.025 0.011 0.51 — b 0.08 0.02 0.40 0.007 0.001 0.022 0.008 0.81 — c 0.08 0.02 0.40 0.007 0.001 0.022 0.023 0.82 — d 0.08 0.10 0.30 0.007 0.001 0.023 0.009 0.89 — e 0.08 0.40 0.20 0.007 0.001 0.024 0.009 0.89 — f 0.11 0.02 0.40 0.007 0.001 0.022 0.010 1.11 — g 0.15 0.02 0.40 0.007 0.001 0.024 0.009 1.51 — h 0.20 0.02 0.40 0.007 0.001 0.022 0.010 2.01 — i 0.26 0.02 0.40 0.007 0.001 0.023 0.009 2.61 — j 0.08 0.02 0.40 0.007 0.001 0.025 0.003 0.80 — k 0.08 0.02 0.40 0.007 0.001 0.025 0.030 0.83 — l 0.31 0.02 0.40 0.007 0.001 0.025 0.008 3.11 — m 0.08 0.60 0.40 0.007 0.001 0.025 0.010 0.81 — n 0.08 0.02 0.15 0.007 0.001 0.025 0.012 0.81 — o 0.08 0.02 1.10 0.007 0.001 0.025 0.011 0.81 — p 0.08 0.02 0.40 0.060 0.001 0.025 0.010 0.81 — q 0.08 0.02 0.40 0.007 0.060 0.025 0.011 0.81 — r 0.08 0.02 0.40 0.007 0.001 0.005 0.012 0.81 — s 0.08 0.02 0.40 0.007 0.001 0.11 0.013 0.81 — t 0.08 0.02 0.40 0.007 0.001 0.025 0.010 0.81 Cr: 0.5, Mb: 0.03 u 0.08 0.02 0.40 0.007 0.001 0.025 0.010 0.81 Cu: 0.06, Ni: 0.15 v 0.08 0.02 0.40 0.007 0.001 0.025 0.009 0.81 Ca: 0.0025, Li: 0.001 w 0.08 0.02 0.40 0.007 0.001 0.025 0.009 0.81 Cr: 0.5, Mb: 0.03 x 0.30 0.02 0.40 0.007 0.001 0.024 0.025 3.03 — (—: not added. Underlined: outside the scope of the present invention.)

TABLE 2 Hot-rolling conditions Rapid Batch annealing conditions Thickness Produc- Kind Heating Final rolling Finish rolling cooling stop Coiling Holding Holding of hot- tion of temperature reduction temperature temperature temperature temperature period rolled sheet No. steel (°C) (%) (° C.) (° C.) (° C.) (° C.) (h) (mm)  1-1* a 1250 15 917 612 575 none* none* 10 1-2 a 1250 15 900 599 519 500 5 10 1-3 a 1250 15 930 601 521 500 1 10 1-4 a 1250 15 921 612 569 500 13 10 1-5 a 1250 15 887 585 577 410 7 10 1-6 a 1250 15 894 621 590 650 5 10  1-7* a 1250 15 886 563 548  350* 5 10  1-8* a 1250 15 898 575 518  730* 5 10  1-9* a 1250 15 895 629 597 500 20*   10  1-10* a 1250 15 886 577 517 400 0.5* 10 2 a 1250 17 909 621 560 500 5 4 3 a 1250 15 895 593 508 500 5 18  4* a  1000*  14*  805*  524*  409* 500 5 10  5* a 1250  14* 907 589 530 500 5 25  6* a 1250  9* 921 623 579 500 5 10 7 b 1250 17 924 611 582 500 5 10 8 c 1250 15 898 572 529 500 5 10 9 d 1250 15 918 621 542 500 5 10 10 e 1250 15 924 597 509 500 5 10 11 f 1250 16 928 568 511 500 5 10 12 g 1250 15 930 639 589 500 5 10 13 h 1250 16 890 627 563 500 5 10 14 i 1250 17 903 594 551 500 5 10 (Underlined: outside the scope of the present invention. *: outside the recommended range.)

TABLE 3 Hot-rolling conditions Final Rapid Thickness Heating rolling Finish cooling Batch annealing conditions of hot- Produc- Kind temper- reduc- rolling stop Coiling Holding Holding rolled tion of ature tion temperature temperature temperature temperature period sheet No. steel (° C.) (%) (° C.) (° C.) (° C.) (° C.) (h) (mm) 15 j 1250 15 907 588 523 500 5 10 16 k 1250 16 914 624 555 500 5 10 17 l 1250 16 924 635 547 500 5 10 18 m 1250 16 895 558 511 500 5 10 19 n 1250 16 910 569 519 500 5 10 20 o 1250 15 914 639 579 500 5 10 21 p 1250 16 913 613 578 500 5 10 22 q 1250 15 927 583 523 500 5 10 23 r 1250 16 915 633 579 500 5 10 24 s 1250 15 893 594 564 500 5 10 25 t 1250 16 918 562 513 500 5 10 26 u 1250 15 913 582 580 500 5 10 27 v 1250 15 901 571 520 500 5 10 28 w 1250 16 908 622 553 500 5 10 29 x 1250 16 924 581 506 500 5 10 (Underlined: outside the scope of the present invention. *: outside the recommended range.)

TABLE 4 Microstructure BF Hardness distribution in thickness average direction Hardness Kind Produc- Solute N Area ratio grain Surface t/4 Central (Hv_(max)- Heavy after Steel of tion amount (%) size portion portion portion Hv_(min))/ cold working No. steel No. (mass %) BF PF P (μm) (Hv) (Hv) (Hv) Hv_(min)(−) workability (Hv) Remarks 1-1 a  1-1* 0.0085 11 86 3 27 176 104 104 0.69 x — Comp. steel 1-2 a 1-2 0.008 12 82 6 31 131 112 113 0.17 ∘ 275 Inventive steel 1-3 a 1-3 0.008 13 85 2 29 141 118 115 0.23 ∘ 283 Inventive steel 1-4 a 1-4 0.007 10 84 6 26 126 115 111 0.14 ∘ 258 Inventive steel 1-5 a 1-5 0.009 14 80 6 32 136 112 107 0.27 ∘ 271 Inventive steel 1-6 a 1-6 0.008 15 81 4 33 124 109 110 0.14 ∘ 255 Inventive steel 1-7 a  1-7* 0.008 13 83 4 41 165 106 107 0.56 x — Comp. steel 1-8 a  1-8* 0.007 14 81 5 45 111 107 106 0.05 ∘ 210 Comp. steel 1-9 a  1-9* 0.009 14 79 7 29 126 113 115 0.12 ∘ 229 Comp. steel 1-10 a  1-10* 0.008 15 82 3 27 163 105 104 0.57 x — Comp. steel (Underlined: outside the scope of the present invention. *: outside the recommended range. BF: bainitic ferrite. PF: polygonal ferrite. P: pearlite. —: not measured because cracking occurred before given rolling reduction during the wedge type compression test. Inventive steel: [(Hv_(max)-Hv_(min))/Hv_(min) ≦ 0.3] and [heavy cold workability = ∘] and [(hardness after working) ≧ 250 Hv]. Comp. steel: the case where any of the requirements for the steel of the present invention is not satisfied.)

TABLE 5 Microstructure BF Hardness distribution in thickness average direction Hardness Kind Produc- Solute N Area ratio grain Surface t/4 Central (Hv_(max) ⁻ Heavy after Steel of tion amount (%) size portion portion portion Hv_(min))/ cold working No. steel No. (mass %) BF PF P (μm) (Hv) (Hv) (Hv) Hv_(min)(−) workability (Hv) Remarks 2 a 2 0.008 27 67 6 15 135 125 127 0.08 ∘ 273 Inventive Steel 3 a 3 0.008 8 88 4 38 126 104 104 0.21 ∘ 251 Inventive steel 4 a   4* 0.003 68 9 23 23 129 106 105 0.23 ∘ 208 Comp. steel 5 a   5* 0.009 4 90 6 55 131 105 104 0.26 ∘ 213 Comp. steel 6 a   6* 0.008 4 88 8 52 127 108 109 0.18 ∘ 205 Comp. steel 7 b 7 0.007 43 52 5 28 128 111 109 0.17 ∘ 292 Inventive steel 8 c 8 0.018 45 48 7 28 134 119 117 0.15 ∘ 296 Inventive steel 9 d 9 0.008 49 47 4 22 132 118 118 0.12 ∘ 308 Inventive steel 10 e 10 0.008 55 42 3 21 142 127 129 0.12 ∘ 318 Inventive steel 11 f 11 0.009 63 32 5 16 144 126 128 0.14 ∘ 311 Inventive steel 12 g 12 0.008 69 20 11 11 151 137 136 0.11 ∘ 325 Inventive steel 13 h 13 0.009 74 11 15 13 162 142 144 0.14 ∘ 331 Inventive steel 14 i 14 0.008 79 6 15 10 166 151 149 0.11 ∘ 351 Inventive steel (Underlined: outside the scope of the present invention. *: outside the recommended range. BF: bainitic ferrite. PF: polygonal ferrite. P: pearlite. —: not measured because cracking occurred before given rolling reduction during the wedge type compression test. Inventive steel: [(Hv_(max)-Hv_(min))/Hv_(min) ≦ 0.3] and [heavy cold workability = ∘] and [(hardness after working) ≧ 250 Hv]. Comp. steel: the case where any of the requirements for the steel of the present invention is not satisfied.)

TABLE 6 Microstructure BF Hardness distribution in thickness average direction Hardness Kind Produc- Solute N Area ratio grain Surface t/4 Central (Hv_(max-) Heavy after Steel of tion amount (%) size portion portion portion Hv_(min))/ cold working No. steel No. (mass %) BF PF P (μm) (Hv) (Hv) (Hv) Hv_(min)(−) workability (Hv) Remarks 15 j 15 0.001 48 49 3 26 129 108 108 0.19 ∘ 181 Comp. steel 16 k 16 0.026 44 51 5 28 129 106 106 0.22 x — Comp. steel 17 l 17 0.007 45 28 27 11 188 153 151 0.25 x — Comp. steel 18 m 18 0.008 58 35 7 27 159 127 125 0.27 x — Comp. steel 19 n 19 0.010 25 71 4 22 119 103 102 0.17 o 221 Comp. steel 20 o 20 0.009 45 52 3 26 131 113 114 0.16 x — Comp. steel 21 p 21 0.009 45 50 5 26 128 106 105 0.22 x — Comp. steel 22 q 22 0.010 47 49 4 29 125 107 107 0.17 x — Comp. steel 23 r 23 0.011 43 55 2 30 122 108 108 0.13 x — Comp. steel 24 s 24 0.012 58 38 4 27 131 108 108 0.21 x — Comp. steel 25 t 25 0.008 55 39 6 17 132 114 116 0.16 ∘ 289 Inventive steel 26 u 26 0.009 58 38 4 16 135 118 118 0.14 ∘ 283 Inventive steel 27 v 27 0.008 50 45 5 19 129 107 107 0.21 ∘ 275 Inventive steel 28 w 28 0.008 59 32 9 17 138 128 127 0.09 ∘ 286 Inventive steel 29 x 29 0.020 68 15 17 11 205 155 154 0.33 x — Comp. steel (Underlined: outside the scope of the present invention. *: outside the recommended range. BF: bainitic ferrite. PF: polygonal ferrite. P: pearlite. —: not measured because cracking occurred before given rolling reduction during the wedge type compression test. Inventive steel: [(Hv_(max)-Hv_(min))/Hv_(min) ≦ 0.3] and [heavy cold workability = ∘] and [(hardness after working) ≧ 250 Hv]. Comp. steel: the case where any of the requirements for the steel of the present invention is not satisfied.)

Tables 4 to 6 show the following. Steels Nos. 1-2 to 1-6, 2, 3, 7 to 14, and 25 to 28 each employed a steel kind which satisfied the requirements regarding the component composition specified in the present invention, and had been produced under the recommended production conditions. As a result, these steels were steels of the present invention which satisfied the requirements regarding the microstructure specified in the present invention, and were on acceptable levels with respect to both heavy cold workability and hardness after working. It was ascertained that hot-rolled steel sheets which show satisfactory heavy cold workability during the working that causes extremely large strains and which show a given hardness (strength) after the working were obtained.

In contrast, steels Nos. 1-1, 1-7 to 1-10, 4 to 6, 15 to 24, and 29 are comparative steels which each do not satisfy at least one of the requirements regarding the component composition and microstructure specified in the present invention. These steels are each not on an acceptable level with respect to heavy cold workability and/or hardness after working.

Specifically, steel No. 1-1 has not undergone batch annealing after the hot rolling, although this steel satisfies the requirements concerning the component composition. This steel has an increased hardness distribution in a thickness direction and is poor at least in heavy cold workability.

Steel No. 1-7, although satisfying the requirements regarding the component composition, has undergone, after the hot rolling, batch annealing in which the holding temperature was too low beyond the recommended range. This steel has an increased hardness distribution in a thickness direction and is poor at least in heavy cold workability.

Meanwhile, steel No. 1-8, although satisfying the requirements regarding the component composition, has undergone, after the hot rolling, batch annealing in which the holding temperature was too high beyond the recommended range. This steel is poor in hardness after working.

Steel No. 1-9, although satisfying the requirements regarding the component composition, has undergone, after the hot rolling, batch annealing in which the holding period was too long beyond the recommended range. This steel is poor in hardness after working.

Meanwhile, steel No. 1-10, although satisfying the requirements regarding the component composition, has undergone, after the hot rolling, batch annealing in which the holding period was too short beyond the recommended range. This steel has an increased hardness distribution in a thickness direction and is poor at least in heavy cold workability.

Steel No 4 satisfies the requirements concerning the component composition, but the heating temperature before the hot rolling was too low beyond the recommended range. This steel has an insufficient solute N amount and is poor in hardness after working.

Steel No. 5, although satisfying the requirements regarding the component composition, has too large a thickness beyond the specified range after the hot rolling. This steel has an insufficient bainitic ferrite content but has too large grain size and is poor in hardness after working.

Steel No. 6, although satisfying the requirements regarding the component composition, has undergone hot rolling in which the final reduction was too low beyond the recommended range. This steel has an insufficient bainitic ferrite content but has too large grain size and is poor in hardness after working.

Steel No. 15 (steel kind j), although produced under the conditions within the recommended ranges, has too low N content. This steel is poor in hardness after working.

Meanwhile, steel No. 16 (steel kind k), although produced under the conditions within the recommended ranges, has too high N content. This steel is poor at least in heavy cold workability.

Steel No. 17 (steel kind 1), although produced under the conditions within the recommended ranges, has too high C content and does not satisfy the requirement 10C+N≦3.0. Pearlite has been excessively formed therein, and this steel is poor at least in heavy cold workability.

Steel No. 18 (steel kind m), although produced under the conditions within the recommended ranges, has too high Si content. This steel is poor at least in heavy cold workability.

Steel No. 19 (steel kind n), although produced under the conditions within the recommended ranges, has too low Mn content. This steel is poor in hardness after working.

Meanwhile, steel No. 20 (steel kind o), although produced under the conditions within the recommended ranges, has too high Mn content. This steel is poor at least in heavy cold workability.

Steel No. 21 (steel kind p), although produced under the conditions within the recommended ranges, has too high P content. This steel is poor at least in heavy cold workability.

Steel No. 22 (steel kind q), although produced under the conditions within the recommended ranges, has too high S content. This steel is poor at least in heavy cold workability.

Steel No. 23 (steel kind r), although produced under the conditions within the recommended ranges, has too low Al content. This steel is poor at least in heavy cold workability.

Meanwhile, steel No. 24 (steel kind s), although produced under the conditions within the recommended ranges except the final reduction during the hot rolling, has too high Al content. This steel is poor at least in heavy cold workability.

Meanwhile, steel No. 29 (steel kind x), although produced under the conditions within the recommended ranges, does not satisfy the requirement 10C+N≦3.0. This steel is poor at least in heavy cold workability.

From these results, the applicability of the present invention was able to be ascertained.

While the present invention has been described in detail and with reference to specific embodiments thereof, it will be apparent to one skilled in the art that various changes and modifications can be made therein without departing from the spirit and scope thereof.

This application is based on a Japanese patent application No. 2014-086747 filed on Apr. 18, 2014, the contents of which are incorporated herein by reference.

INDUSTRIAL APPLICABILITY

The hot-rolled steel sheet of the present invention shows satisfactory workability in cold working and shows given hardness after the working. This hot-rolled steel sheet is useful as a steel material for use in producing, in particular, various automotive parts, such as, for example, transmission parts, e.g., gears, and cases. 

1. A hot-rolled steel sheet excellent in heavy cold workability and surface hardness after working, wherein: a sheet thickness is from 3 to 20 mm; a component composition comprises, in mass %, C: more than 0% and 0.3% or less, Si: more than 0% and 0.5% or less, Mn: from 0.2 to 1%, P: more than 0% and 0.05% or less, S: more than 0% and 0.05% or less, Al: from 0.01 to 0.1%, and N: from 0.008 to 0.025%, with the remainder being iron and unavoidable impurities, wherein solute N: 0.007% or more and the contents of C and N satisfy the relationship 10C+N≦3.0; a microstructure comprises, in terms of area ratio relative to an entire microstructure, bainitic ferrite: 5% or more, pearlite: less than 20%, and remainder: polygonal ferrite; an average grain size of the bainitic ferrite is in a range of from 3 to 50 μm; and (Hv_(max)−Hv_(min))/Hv_(min)≦0.3, wherein, in a hardness distribution in a thickness direction, Hv_(max) and Hv_(min)≦0.3, are respectively the maximal value and the minimal value of Vickers hardness values of three portions that are a surface portion, a t/4 portion where t is the sheet thickness, and a central portion.
 2. The hot-rolled steel sheet according to claim 1, wherein the component composition further comprises at least one member of the following (a) to (e): (a) at least one member selected from the group consisting of Cr: more than 0% and 2% or less and Mo: more than 0% and 2% or less, in mass %; (b) at least one member selected from the group consisting of Ti: more than 0% and 0.2% or less, Nb: more than 0% and 0.2% or less, and V: more than 0% and 0.2% or less, in mass %; (c) B: more than 0% and 0.005% or less in mass %; (d) at least one member selected from the group consisting of Cu: more than 0% and 5% or less, Ni: more than 0% and 5% or less, and Co: more than 0% and 5% or less, in mass %; and (e) at least one member selected from the group consisting of Ca: more than 0% and 0.05% or less, REM: more than 0% and 0.05% or less, Mg: more than 0% and 0.02% or less, Li: more than 0% and 0.02% or less, Pb: more than 0% and 0.5% or less, and Bi: more than 0% and 0.5% or less. 